A nickel-based alloy

ABSTRACT

A nickel-based alloy composition consisting, in weight percent, of: between 3.5 and 6.5% chromium, between 0.0 and 12.0% cobalt, between 4.5 and 11.5% tungsten, between 0.0 and 0.5% molybdenum, between 3.5 and 7.0% rhenium, between 1.0 and 3.7% ruthenium, between 3.7 and 6.8% aluminium, between 5.0 and 9.0% tantalum, between 0.0 and 0.5% hafnium, between 0.0 and 0.5% niobium, between 0.0 and 0.5% titanium, between 0.0 and 0.5% vanadium, between 0.0 and 0.1% silicon, between 0.0 and 0.1% yttrium, between 0.0 and 0.1% lanthanum, between 0.0 and 0.1% cerium, between 0.0 and 0.003% sulphur, between 0.0 and 0.05% manganese, between 0.0 and 0.05% zirconium, between 0.0 and 0.005% boron, between 0.0 and 0.01% carbon, the balance being nickel and incidental impurities.

The present invention relates to a nickel-based single crystalsuperalloy composition designed for high performance jet propulsionapplications. The alloy—a fourth generation single crystal nickel-basedsuperalloy—exhibits a combination of creep resistance and oxidationresistance which is comparable or better than equivalent grades ofalloy. The density, cost, processing and long term stability of thealloy have also been considered in the design of the new alloy.

Examples of typical compositions of fourth generation nickel-basedsingle crystal superalloys are listed in Table 1. These alloys may beused for the manufacture of rotating/stationary turbine blades used ingas turbine engines.

TABLE 1 Nominal composition in wt. % of commercially used fourthgeneration single crystal turbine blade alloys. Alloy Al Cr Co Mo Re RuW Ti Ta Hf PW1497 5.6 2.0 16.5 2.0 6.0 3.0 6.0 0.0 8.3 0.2 TMS-162 5.82.9 5.8 3.9 4.9 6.0 5.8 0.0 5.6 0.1 TMS-138A 5.7 3.2 5.8 2.8 5.8 3.6 5.60.0 5.6 0.0

It is an aim of the invention is to provide an alloy which has similaror improved high temperature behaviour in comparison to the fourthgeneration alloys listed in Table 1.

The present invention provides a nickel-based alloy compositionconsisting, in weight percent, of: between 3.5 and 6.5% chromium,between 0.0 and 12.0% cobalt, between 4.5 and 11.5% tungsten, between0.0 and 0.5% molybdenum, between 3.5 and 7.0% rhenium, between 1.0 and3.7% ruthenium, between 3.7 and 6.8% aluminium, between 5.0 and 9.0%tantalum, between 0.0 and 0.5% hafnium, between 0.0 and 0.5% niobium,between 0.0 and 0.5% titanium, between 0.0 and 0.5% vanadium, between0.0 and 0.1% silicon, between 0.0 and 0.1% yttrium, between 0.0 and 0.1%lanthanum, between 0.0 and 0.1% cerium, between 0.0 and 0.003% sulphur,between 0.0 and 0.05% manganese, between 0.0 and 0.05% zirconium,between 0.0 and 0.005% boron, between 0.0 and 0.01% carbon, the balancebeing nickel and incidental impurities. This composition provides a goodbalance between cost, density and creep and oxidation resistance.

In an embodiment, the nickel-based alloy composition consists, in weightpercent, of between 4.0 and 5.0% chromium. Such an alloy is particularlyresistant to TCP formation whilst still having good oxidationresistance.

In an embodiment, the nickel-based alloy composition contains at least0.1 wt % cobalt to increase creep resistance and lower the γ′ solvustemperature thereby to increase the solutioning window.

In an embodiment, the nickel-based alloy composition consists, in weightpercent, of between 7.0 and 11.0% cobalt. Such an alloy has improvedresistance to creep deformation with a limited level of creep anisotropy(orientation dependence) being observed and has increased ease ofprocessing due to a reduced γ′ solvus temperature. A maximum amount ofcobalt of 10.9% further limits creep anisotropy.

In an embodiment, the nickel-based alloy composition consists, in weightpercent, of between 7.0 and 9.5% tungsten. This composition strikes acompromise between reduced cost, low weight and creep resistance.

In an embodiment, the nickel-based alloy composition consists, in wt %,of at least 7.1% tungsten in order to achieve high creep resistance.

In an embodiment, the nickel-based alloy composition consists, in weightpercent, of between 3.7 and 6.6% aluminium, preferably between 5.1 and6.6% aluminium, or more preferably 5.5 and 6.6% aluminium. Thiscomposition achieves high creep resistance and reduced density alongsideincreased oxidation resistance.

In an embodiment, the nickel-based alloy composition consists, in weightpercent, of between 5.0 and 9.0% tantalum. This provides a balancebetween creep resistance, ease of manufacture (based upon solutioningwindow) and density and/or prevents the possibility of formation of theEta (ε) phase Ni₃Ta. Preferably the alloy consists of between 5.0 and7.3% tantalum. This reduces the cost and density of the alloy further,increases the solutioning window as well as the propensity for ε phaseformation.

In an embodiment, the nickel-based alloy composition consists, in weightpercent, of 0.1% or more molybdenum. This is advantageous for improvedcreep resistance.

In an embodiment, the nickel-based alloy composition consists, in weightpercent of, between 4.5 and 6.0% rhenium, or more preferably 5.3 and6.0% rhenium. This composition provides a good balance of creepresistance, density, resistance to TCP formation and cost.

In an embodiment, the nickel-based alloy composition consists, in weightpercent of, between 2.0 and 3.0% ruthenium. This composition provides agood balance of creep resistance and cost. An even better compromise isprovided in the range of 2.1 and 2.9% ruthenium.

In an embodiment, the nickel-based alloy composition consists, in weightpercent, of between 0.0 and 0.2% hafnium. This is optimum for tying upincidental impurities in the alloy, for example, carbon.

In an embodiment, the nickel-based alloy composition is such that thefollowing equation is satisfied in which W_(Ta) and W_(Al) are theweight percent of tantalum and aluminium in the alloy respectively33≤W_(Ta)+5.1 W_(Al)≤39. This is advantageous as it allows a suitablevolume fraction γ′ to be present.

In an embodiment, the nickel-based alloy composition is such that thefollowing equation is satisfied in which W_(Ta) and W_(Al) are theweight percent of tantalum and aluminium in the alloy respectively2.2≤5.15 W_(Al)−0.5 W_(Al) ²−W_(Ta); preferably 2.9≤5.15 W_(Al)−0.5W_(Al) ²−W_(Ta). This is advantageous as it allows a suitablesolutioning window for the alloy to allow for heat-treatment processes.

In an embodiment, the nickel-based alloy composition is such that thefollowing equation is satisfied in which W_(Ru) and W_(Re) are theweight percent of ruthenium and rhenium in the alloy respectively4.5≥W_(Ru)+0.225 W_(Re); preferably 3.9≥W_(Ru)+0.225 W_(Re). This isadvantageous as it results in an alloy with a relatively low cost.

In an embodiment, the nickel-based alloy composition is such that thefollowing equation is satisfied in which W_(Re) and W_(W) are the weightpercent of rhenium and tungsten in the alloy respectively 15.8≥1.13W_(Re)+W_(W); preferably 14.4≥1.13 W_(Re)+W_(W). This is advantageous asit results in an alloy with a relatively low density.

In an embodiment, the nickel-based alloy composition is such that thefollowing equation is satisfied in which W_(Re), W_(Mo) and W_(W) arethe weight percent of rhenium, molybdenum and tungsten in the alloyrespectively 21.9≤2.92 W_(Re)+(W_(W)+W_(Mo)); preferably 24.6≤2.92W_(Re)+(W_(W)+W_(Mo)). This is advantageous as it results in an alloywith a high creep resistance.

In an embodiment, in the nickel-based alloy composition, the sum of theelements niobium, titanium and vanadium, in weight percent, is less than1%, preferably 0.5% or less. This means that those elements do not havetoo much of a deleterious effect on environmental resistance of thealloy.

In an embodiment, in the nickel-based alloy composition, the sum of theelements niobium, titanium, vanadium and tantalum is between 5.0-9.0 wt.%, preferably 5.0-7.3 wt. %. This results in a preferred volume fractionof γ′ and APB energy.

In an embodiment, the nickel-based alloy composition has between 60 and70% volume fraction γ′.

In an embodiment, a single crystal article is provided, formed of thenickel-based alloy composition of any of the previous embodiments.

In an embodiment, a turbine blade for a gas turbine engine is provided,formed of an alloy according to any of the previous embodiments.

In an embodiment, a gas turbine engine comprising the turbine blade ofthe previous embodiment is provided.

The term “consisting of” is used herein to indicate that 100% of thecomposition is being referred to and the presence of additionalcomponents is excluded so that percentages add up to 100%.

The invention will be more fully described, by way of example only, withreference to the accompanying drawings in which:

FIG. 1 shows the partitioning coefficient for the main components in thealloy design space;

FIG. 2 is a contour plot showing the effect of γ′ forming elementsaluminium and tantalum on volume fraction of γ′ for alloys within thealloy design space, determined from phase equilibrium calculationsconducted at 900° C.;

FIG. 3 is a contour plot showing the effect of elements aluminium andtantalum on anti-phase boundary energy, for alloys with a volumefraction of γ′ between 60-70% at 900° C.;

FIG. 4 is a contour plot showing the effect of elements aluminium andtantalum on the solutioning window for alloys with a volume fraction ofγ′ between 60-70% at 900° C.;

FIG. 5 is a contour plot showing the effect of rhenium and rutheniumcontent on raw elemental cost, for alloys with a volume fraction of γ′between 60-70% at 900° C. with tantalum between 5-9 wt. %;

FIG. 6 is a contour plot showing the effect rhenium and tungsten ondensity, for alloys with a volume fraction of γ′ between 60-70% at 900°C. with tantalum between 5-9 wt. %;

FIG. 7 is a contour plot showing the effect of elements rhenium andtungsten on the creep resistance, for alloys with a volume fraction ofγ′ between 60-70% at 900° C. with tantalum between 5-9 wt. %, whichcontain 0 wt. % ruthenium;

FIG. 8 is a contour plot showing the effect of elements rhenium andtungsten on the creep resistance, for alloys with a volume fraction ofγ′ between 60-70% at 900° C. with tantalum between 5-9 wt. %, whichcontain 1 wt. % ruthenium;

FIG. 9 is a contour plot showing the effect of elements rhenium andtungsten on the creep resistance, for alloys with a volume fraction ofγ′ between 60-70% at 900° C. with tantalum between 5-9 wt. %, whichcontain 2 wt. % ruthenium;

FIG. 10 is a contour plot showing the effect of elements rhenium andtungsten on the creep resistance, for alloys with a volume fraction ofγ′ between 60-70% at 900° C. with tantalum between 5-9 wt. %, whichcontain 3 wt. % ruthenium;

FIG. 11 is a contour plot showing the effect of elements chromium andtungsten on microstructural stability, for alloys with a volume fractionof γ′ between 60-70% at 900° C. with tantalum between 5-9 wt. % andbetween 1-3 wt. % ruthenium, which contain 4 wt. % rhenium;

FIG. 12 is a contour plot showing the effect of elements chromium andtungsten on microstructural stability, for alloys with a volume fractionof γ′ between 60-70% at 900° C. with tantalum between 5-9 wt. % andbetween 1-3 wt. % ruthenium, which contain 5 wt. % rhenium;

FIG. 13 is a contour plot showing the effect of elements chromium andtungsten on microstructural stability, for alloys with a volume fractionof γ′ between 60-70% at 900° C. with tantalum between 5-9 wt. % andbetween 1-3 wt. % ruthenium, which contain 6 wt. % rhenium;

FIG. 14 is a contour plot showing the effect of elements chromium andtungsten on microstructural stability, for alloys with a volume fractionof γ′ between 60-70% at 900° C. with tantalum between 5-9 wt. % andbetween 1-3 wt. % ruthenium, which contain 7 wt. % rhenium;

FIG. 15 is a contour plot showing the effect of cobalt on the γ′ solvustemperature for alloys with different ratios of aluminium to tantalum,where the alloys have a volume fraction of γ′ between 60-70% at 900° C.with tantalum between 5-9 wt. %;

FIG. 16 shows the time to 1% creep strain for alloy ABD-2 of the presentinvention (circles) compared with the fourth generation single crystalturbine blade alloy TMS-138A (triangles);

FIG. 17 shows the time to rupture for alloy ABD-2 of the presentinvention (circles) compared with the fourth generation single crystalturbine blade alloy TMS-138A (triangles); and

FIG. 18 is a plot of measured weight change for the fourth generationsingle crystal turbine blade alloy TMS-138A (triangles) and alloy ABD-2of the present invention (circles) when oxidised in air at 1000° C.

Traditionally, nickel-based superalloys have been designed throughempiricism. Thus their chemical compositions have been isolated usingtime consuming and expensive experimental development, involvingsmall-scale processing of limited quantities of material and subsequentcharacterisation of their behaviour. The alloy composition adopted isthen the one found to display the best, or most desirable, combinationof properties. The large number of possible alloying elements indicatesthat these alloys are not entirely optimised and that improved alloysare likely to exist.

In superalloys, generally additions of chromium (Cr) and aluminium (Al)are added to impart resistance to oxidation, cobalt (Co) is added toimprove resistance to sulphidisation. For creep resistance, molybdenum(Mo), tungsten (W), Co, rhenium (Re) and sometimes ruthenium (Ru) areintroduced, because these retard the thermally-activated processes—suchas, dislocation climb—which determine the rate of creep deformation. Topromote static and cyclic strength, aluminium (Al), tantalum (Ta) andtitanium (Ti) are introduced as these promote the formation of theprecipitate hardening phase gamma-prime (γ′). This precipitate phase iscoherent with the face-centered cubic (FCC) matrix phase which isreferred to as gamma (γ).

A modelling-based approach used for the isolation of new grades ofnickel-based superalloys is described here, termed the“Alloys-By-Design” (ABD) method. This approach utilises a framework ofcomputational materials models to estimate design relevant propertiesacross a very broad compositional space. In principle, this alloy designtool allows the so called inverse problem to be solved; identifyingoptimum alloy compositions that best satisfy a specified set of designconstraints.

The first step in the design process is the definition of an elementallist along with the associated upper and lower compositional limits. Thecompositional limits for each of the elemental additions considered inthis invention—referred to as the “alloy design space”—are detailed inTable 2.

TABLE 2 Alloys design space in wt. % searched using the“Alloys-by-Design” method. Cr Co Re W Al Ta Ru Min 2.0 0.0 3.0 4.0 3.03.0 0.0 Max 10.0 16.0 10.0 12.0 9.0 16.0 4.0

The second step relies upon thermodynamic calculations used to calculatethe phase diagram and thermodynamic properties for a specific alloycomposition. Often this is referred to as the CALPHAD method (CALculatePHAse Diagram). These calculations are conducted at the servicetemperature for the new alloy (900° C.), providing information about thephase equilibrium (microstructure).

A third stage involves isolating alloy compositions which have thedesired microstructural architecture. In the case of single crystalsuperalloys which require superior resistance to creep deformation, thecreep rupture life is maximised when the volume fraction of theprecipitate hardening phase γ′ lies between 60%-70%. It is alsonecessary that the γ/γ′ lattice misfit should conform to a small value,either positive or negative, since coherency is otherwise lost; thuslimits are placed on its magnitude. The lattice misfit 6 is defined asthe mismatch between γ and γ′ phases, and is determined according to

$\begin{matrix}{\delta = \frac{2\left( {a_{\gamma^{\prime}} - a_{\gamma}} \right)}{a_{\gamma^{\prime}} + a_{\gamma}}} & (1)\end{matrix}$

where a_(γ) and a_(γ′) are the lattice parameters of the γ and γ′phases.

Rejection of alloy on the basis of unsuitable microstructuralarchitecture is also made from estimates of susceptibility totopologically close-packed (TCP) phases. The present calculationspredict the formation of the deleterious TCP phases sigma (σ), P and mu(μ) using CALPHAD modelling.

Thus the model isolates all compositions in the design space which arecalculated to result in a volume fraction of γ′ of between 60 and 70%,which have a lattice misfit γ′ of less than a predetermined magnitudeand have a total volume fraction of TCP phases below a predeterminedmagnitude.

In the fourth stage, merit indices are estimated for the remainingisolated alloy compositions in the dataset. Examples of these include:creep-merit index (which describes an alloy's creep resistance basedsolely on mean composition), anti-phase boundary (APB) energy, density,cost and solutioning window.

In the fifth stage, the calculated merit indices are compared withlimits for required behaviour, these design constraints are consideredto be the boundary conditions to the problem. All compositions which donot fulfil the boundary conditions are excluded. At this stage, thetrial dataset will be reduced in size quite markedly.

The final, sixth stage involves analysing the dataset of remainingcompositions. This can be done in various ways. One can sort through thedatabase for alloys which exhibit maximal values of the meritindices—the lightest, the most creep resistant, the most oxidationresistant, and the cheapest for example. Or alternatively, one can usethe database to determine the relative trade-offs in performance whicharise from different combination of properties.

The example five merit indices are now described.

The first merit index is the creep-merit index. The overarchingobservation is that time-dependent deformation (i.e. creep) of a singlecrystal superalloy occurs by dislocation creep with the initial activitybeing restricted to the γ phase. Thus, because the fraction of the γ′phase is large, dislocation segments rapidly become pinned at the γ/γ′interfaces. The rate-controlling step is then the escape of trappedconfigurations of dislocations from γ/γ′ interfaces, and it is thedependence of this on local chemistry which gives rise to a significantinfluence of alloy composition on creep properties.

A physically-based microstructure model can be invoked for the rate ofaccumulation of creep strains {dot over (ε)} when loading is uniaxialand along the

001

crystallographic direction. The equation set is

$\begin{matrix}{{{\overset{.}{ɛ}}_{\langle 001\rangle} = {\frac{16}{\sqrt{6}}\rho_{m}\varphi_{p}{D_{eff}\left( {1 - \varphi_{p}} \right)}\left( {{1/\varphi_{p}^{1/3}} - 1} \right)\sin \; h\left\{ \frac{\sigma \; b^{2}\omega}{\sqrt{6}K_{CF}{kT}} \right\}}}} & (2) \\{{\overset{.}{\rho}}_{m} = {C{\overset{.}{ɛ}}_{\langle 001\rangle}}} & (3)\end{matrix}$

where ρ_(m) is the mobile dislocation density, ϕ_(p) is the volumefraction of the γ′ phase, and ω is width of the matrix channels. Theterms σ and T are the applied stress and temperature, respectively. Theterms b and k are the Burgers vector and Boltzmann constant,respectively. The term K_(CF)=1+2ϕ_(p) ^(1/3)3√{square root over(3π)}(1−ϕ_(p) ^(1/3)) is a constraint factor, which accounts for theclose proximity of the cuboidal particles in these alloys. Equation 3describes the dislocation multiplication process which needs an estimateof the multiplication parameter C and the initial dislocation density.The term D_(eff) is the effective diffusivity controlling the climbprocesses at the particle/matrix interfaces.

Note that in the above, the composition dependence arises from the twoterms ϕ_(p) and D_(eff). Thus, provided that the microstructuralarchitecture is assumed constant (microstructural architecture is mostlycontrolled by heat treatment) so that ϕ_(p) is fixed, any dependenceupon chemical composition arises through D_(eff). For the purposes ofthe alloy design modelling described here, it turns out to beunnecessary to implement a full integration of Equations 2 and 3 foreach prototype alloy composition. Instead, a first order merit indexM_(creep) is employed which needs to be maximised, which is given by

$\begin{matrix}{M_{creep} = {\sum\limits_{i}\; {x_{i}/{\overset{\sim}{D}}_{i}}}} & (4)\end{matrix}$

where x_(i) is the atomic fraction of solute i in the alloy and {tildeover (D)}_(i) is the appropriate interdiffusion coefficient.

The second merit index is for anti-phase boundary (APB) energy. Faultenergies in the γ′ phase—for example, the APB energy—have a significantinfluence on the deformation behaviour of nickel-based superalloys.Increasing the APB energy has been found to improve mechanicalproperties including, tensile strength and resistance to creepdeformation. The APB energy was studied for a number of Ni—Al—X systemsusing density functional theory. From this work the effect of ternaryelements on the APB energy of the γ′ phase was calculated, linearsuperposition of the effect for each ternary addition was assumed whenconsidering complex multicomponent systems, resulting in the followingequation,

γ_(APB)=195−1.7x _(Mo)+1.7x _(Mo)+4.6x _(W)+27.1x _(Ta)+21.4x _(Nb)+15x_(Ti)  (5)

where, x_(Cr), x_(Mo), x_(W), x_(Ta), x_(Nb) and x_(Ti) represent theconcentrations, in atomic percent, of Cr, Mo, W, Ta, Nb and Ti in the γ′phase, respectively. The composition of the γ′ phase is determined fromphase equilibrium calculations.

The third merit index is density. The density, ρ, was calculated using asimple rule of mixtures and a correctional factor, where, ρ_(t) is thedensity for a given element and x_(i) is the atomic fraction of thealloy element.

ρ=1.05[Σ_(i) x _(i)ρ_(i)]  (6)

The fourth merit index was cost. In order to estimate the cost of eachalloy a simple rule of mixtures was applied, where the weight fractionof the alloy element, x_(i), was multiplied by the current (2015) rawmaterial cost for the alloying element, c_(i).

Cost=Σ_(i) x _(i) c _(i)  (7)

The estimates assume that processing costs are identical for all alloys,i.e. that the product yield is not affected by composition.

A fifth merit index is the solutioning window. By conductingthermodynamic modelling (CALPHAD) calculations across a range oftemperatures the solutioning window for each alloy can be calculated.This value—measured in degrees Celsius—can be used to determine if agiven alloy is amenable to conventional manufacturing processes used forthe production of single crystal turbine blades. Typically thesolutioning window must be greater than 50° C. to allow for a solutionheat treatment. The solution heat treatment is conducted in the singlephase region, at this point the alloy will reside solely within the γphase field. This solution heat treatment is necessary to homogenise thecomposition of the as cast alloy which may be highly segregated. Inorder determine the solution heat treatment window the phaseequilibrium—or more specifically phase transformations—must bedetermined over a temperature range. The temperature at which completeddissolution of the γ′ phase (known as the γ′ solvus temperature) occursmust be known, as must the solidus temperature. The difference betweenthe solidus temperature and the γ′ solvus temperature will give thesolutioning window. So the solutioning window index calculates as thedifference between the solidus temperature and the γ′ solvustemperature.

The ABD method described above was used to isolate the inventive alloycomposition. The design intent for this alloy was to isolate thecomposition of a fourth generation single crystal nickel-basedsuperalloy that exhibits a combination of creep resistance and oxidationresistance which is comparable or better than equivalent grades ofalloy. The density, cost, processing and long term stability of thealloy have also been considered in the design of the new alloy.

The material properties—determined using the ABD method—for thecommercially used fourth generation single crystal turbine blade alloysare is listed in Table 3. The design of the new alloy was considered inrelation to the predicted properties listed for these alloys. Thecalculated material properties for an alloy ABD-2 with a nominalcomposition according to Table 4 and in accordance with the presentinvention are also given.

TABLE 3 Calculated phase fractions and merit indices made with the“Alloys-by-Design” software. Results for fourth generation singlecrystal turbine blades listed in Table 1 and the nominal composition ofthe new alloy ABD-2 listed in Table 4. Creep Merit Index γ/γ′Solutioning Phase Fractions (m⁻² s × Density Cost γ_(APB(111)) MisfitWindow Alloy γ′ P μ P + μ 10⁻¹⁵) (g/cm³) ($/lb) (mJ/m²) (%) (° C.)PW1497 0.61 0.000 0.054 0.054 14.6 9.2 296 322 −0.02 43 TMS-162 0.640.000 0.060 0.060 13.8 9.0 467 274 −0.44 94 TMS-138A 0.60 0.016 0.0320.048 12.3 9.0 328 278 −0.19 116 ABD-2 0.66 0.000 0.037 0.037 12.3 8.9260 274 −0.14 62

Optimisation of the alloy's microstructure—primarily comprised of anaustenitic face centre cubic (FCC) gamma phase (γ) and the ordered L1₂precipitate phase (γ′)—was required to maximise creep resistance. Avolume fraction of the γ′ phase between 60-70% is generally regarded asoptimum as this microstructure is known to provide the maximum level ofcreep resistance in single crystal blade alloys. A volume fraction γ′ ofbetween 60 and 70% was the target for the present alloy but theinventive alloy may deviate from this target.

The partitioning coefficient for each element included in the alloydesign space was determined from phase equilibrium calculationsconducted at 900° C., FIG. 1. A partitioning coefficient of unitydescribes an element with equal preference to partition to the γ or γ′phase. A partitioning coefficient less than unity describes an elementwhich has a preference for the γ′ phase, the closer the value to zerothe stronger the preference. The greater the value above unity the morean element prefers to reside within the γ phase. The partitioningcoefficients for aluminium and tantalum show that these are strong γ′forming elements. The elements chromium, cobalt, rhenium, ruthenium andtungsten partition preferably to the γ phase. For the elementsconsidered within the alloy design space, aluminium and tantalumpartition most strongly to the γ′ phase. Hence, aluminium and tantalumlevels were controlled to produce the desired γ′ volume fraction.

FIG. 2 shows the effect the elements which are added to form the γ′phase—predominantly aluminium and tantalum—have on the fraction of γ′phase in the alloy at the operation temperature, 900° C. in thisinstance. For the design of this alloy compositions which result in avolume fraction of γ′ between 60-70% were considered. Hence between 3.7and 6.8 weight percent (wt. %) of aluminium was required.

The change in γ′ volume fraction was related to the change in aluminiumand tantalum content according to the formula

f(γ′)=W _(Ta)+5.1W _(Al)

where, f(γ′) is a numerical value which ranges between 33 and 39 for analloy with the desired γ′ fraction, between 0.6 and 0.7 in this case,and W_(Ta) and W_(Al) are the weight percent of tantalum and aluminiumin the alloy, respectively.

Optimisation of aluminium and tantalum levels was also required toincrease the anti-phase boundary (APB) energy of the γ′ phase. The APBenergy is strongly dependent upon the chemistry of the γ′ phase. FIG. 3shows the influence of aluminium and tantalum on the APB energy.Compositions where the APB energy was equivalent to or greater thancurrent fourth generation single crystal alloys (˜270 mJ/m²) weredetermined. Modelling calculation showed that tantalum levels in thealloy greater than 5.0 wt. % (at Al=3.7-6.8 wt. %, meaning 60-70% volumeγ′) produce an alloy with an acceptably high APB energy and so highcreep resistance and at the same time producing a high enough volumefraction of γ′.

With the minimum tantalum concentration, aluminium additions beinglimited to a maximum of 6.6 wt. % is desirable so that the desirable γ′volume fraction of 60-70% is achieved, FIG. 2. Therefore, Alconcentration of between 3.7 and 6.6 wt. % is desirable to achieve boththe desired γ′ volume fraction and a high APB energy. The maximumtantalum content will be explained below with reference to FIG. 4 andresults in a tantalum range of 5.0-9.0 wt. % and a preferred range of5.0 to 7.3 wt % this results from the preferred combination of APBenergy and solutioning window (dealt with below). That is, the preferredminimum levels of tantalum ensure a higher APB energy for any givenamount of aluminium and a level of at least 270 mJ/m² in the range ofaluminium for the alloy. From FIG. 2 it is seen that in particular forhigher lower levels of tantalum, concentrations of aluminium of 5.1 wt.% or more, preferably 5.5 wt. % or more produce the desired volumefraction of γ′. Therefore, it is preferable to have the ratio ofaluminium to tantalum, in weight percent, ranging between 0.41 (Al=3.7wt. %, Ta=9.0 wt. %) and 1.36 (Al=6.8 wt. %, Ta=5.0 wt. %), or morepreferably ranging between 0.70 (Al=5.1 wt. %, Ta=7.3 wt. %) and 1.32(Al=6.6 wt. %, Ta=5.0 wt. %), or even more preferably between 0.75(Al=5.5 wt. %, Ta=7.3 wt. %) and 1.32 (Al=6.6 wt. %, Ta=5.0 wt. %).

Niobium, titanium, vanadium elements behave in a similar way to that oftantalum i.e. they are gamma prime forming elements which increaseanti-phase boundary energy. These elements can optionally be added tothe alloy. The benefits of this may include lower cost and density incomparison to tantalum. However, additions of these elements must belimited as they can have a negative impact on the environmentalresistance of the alloy. Therefore, those elements can each be presentin an amount of up to 0.5 wt. %. Preferably those elements aresubstituted for tantalum meaning that the sum of the elements consistingof niobium, titanium, vanadium and tantalum is preferably limited to5.0-9.0 wt. %, more preferably 5.0-7.3 wt. % which are the preferredranges for tantalum. Independently, in an embodiment, the sum of theelements consisting of niobium, titanium and vanadium is preferablylimited to below 1.0 wt. % and preferably below 0.5 wt. % so as to avoidreduction in environmental resistance of the alloy.

The balance of aluminium and tantalum can be adjusted such that there isa balance between desired target volume fraction of γ′ as well as asufficiently high APB energy. However, consideration must also be givento the processing of the alloy. One such consideration is thesolutioning window; there should exist a sufficient temperature rangewindow, below the melting temperature of the alloy, across which onlythe γ phase is stable. As the solutioning window depends upon thedissolution of the γ′ phase it is strongly influenced by γ′ chemistry,hence, aluminium and tantalum content. This solutioning heat treatmentis used to remove any residual microsegregation and eutectic mixturesrich in γ′ which might occur during the casting processes used toproduce the single crystal alloy. It is preferred that the solutioningwindow is greater than 50° C. to allow for conventional processingmethods. FIG. 4 shows the solutioning window magnitude (in ° C.) forvarying wt % Al and Ta with a volume fraction γ′ of 60-70%. From thisFigure it can be seen that limiting the tantalum content to 9.0 wt. %ensures that the alloy has a suitable solutioning window. Preferably thetantalum content is limited to 7.3 wt. % as this produces an alloy witha solutioning window greater than 60° C. further improving theprocessing of the alloy.

The change in solutioning window was related to the change in aluminiumand tantalum content according to the formula

f(T _(sol.))=5.15W _(Al)−0.57W _(Al) ² −W _(Ta)

where, f(T_(sol.)) is a numerical value which is greater than 2.2 toproduce an alloy with a solutioning window greater than or equal to 50°C. f(T_(sol.)) is preferably greater than 2.9 to produce and alloy witha solutioning window greater than 60° C.

For the alloys which satisfied the previously described requirements(volume fraction of γ′ between 60-70%, APB energy greater than 270mJ/m², solutioning window greater than 50° C.) the levels of refractoryelements were determined for creep resistance and oxidation performance.For fourth generation single crystal turbine blades additions of theelements ruthenium, rhenium and tungsten are made to impart substantialcreep performance, this is described later with reference to FIGS. 7-10.However, the elements rhenium and ruthenium strongly effect cost, FIG.5. The elements tungsten and rhenium significantly increase alloydensity, FIG. 6. Moreover, elements such as rhenium, tungsten andchromium (chromium is added for oxidation resistance) must be suitablybalanced such that a balance between creep resistance and oxidation isachieved without resulting in a microstructurally unstable alloy whichis prone to the formation of deleterious TCP phases, FIGS. 11-14. Thus,a complex balance between trade-offs in cost, density, creep resistance,oxidation resistance and microstructural stability must be managed, theprocess for optimising these trade-offs is described below withreference to FIGS. 5-14.

The current (2015) raw material cost for the elements ruthenium andrhenium is substantial. Therefore, to optimise the design of the alloylevels of ruthenium and rhenium are selected which best manage thetrade-off between the cost and creep resistance in the presentinvention. In FIG. 5 a contour plot shows the effect which levels ofruthenium and rhenium have on alloy cost for an alloy of 60-70% γ′volume fraction at 900° C. It is seen that ruthenium has the strongestinfluence on alloy cost. Thus, the ruthenium content in the alloy islimited to 3.7 wt. % ensuring that the cost of the present invention isequivalent to or less than current grades of fourth generation alloy. Itis preferred that the ruthenium content is limited to 3.0 wt. % or lessto ensure an optimal balance between cost and creep resistance.

In order to limit the cost of the alloy, additions of ruthenium andrhenium preferably adhere to the following Equation,

f(Cost)=W _(Ru)+0.225W _(Re)

where, f(Cost) is a numerical value which is less than or equal to 4.5to produce an alloy with a cost of 300$/lb or less and W_(Ru) and W_(Re)is the weight percent of ruthenium and rhenium in the alloyrespectively. Preferably the numerical value for f(Cost) is less than orequal to 3.9 as this produces an alloy with a lower cost of 260$/lb orless.

The additions of the elements tungsten, rhenium and ruthenium areoptimised in order to design an alloy which is highly resistant to creepdeformation. The creep resistance was determined by using the creepmerit index model. It is desirable to maximise the creep merit index asthis is associated with an improved creep resistance. The influencewhich tungsten, rhenium and ruthenium have on creep resistance ispresented in FIGS. 7-10. It is seen that increasing the levels oftungsten, rhenium and ruthenium improve creep resistance. However, thequantities of tungsten and rhenium required mean that they have a stronginfluence on alloy density, FIG. 6. The calculations to produce thegraphs of FIGS. 6-10 are done such that the γ′ volume fraction at 900°C. is between 60 and 70%. Therefore the trade-off between creepresistance and alloy density must be balanced.

In order to limit the density of the alloy additions of tungsten andrhenium preferably adhere to the following Equation,

f(Desnity)=1.13W _(Re) +W _(W)

where, f(Density) is a numerical value which is less than or equal to15.8 to produce an alloy with a density of 9.0 g/cm³ or less and W_(W)is the weight percent of tungsten in the alloy. Preferably the numericalvalue for f(Density) is less than or equal to 14.4 as this produces analloy with a density of 8.9 g/cm³ or less.

Current fourth generation single crystal alloys have a creep merit indexof 12×10⁻¹⁵ m⁻²s or greater (see Table 3). This level of creepresistance is desirably attained in combination with a low density ofless than 9.0 g/cm³ or preferably 8.9 g/cm³. In FIGS. 7-10 the contoursfrom FIG. 6 (dashed lines) which show the effect which rhenium andtungsten have on density are superimposed on the effect which rhenium,tungsten and ruthenium have on creep merit index.

In order to attain a minimum creep merit index of 12×10⁻¹⁵ m⁻²s, thealloy contains at least 1.0 wt. % of ruthenium. Preferably the rutheniumcontent is 2.0 wt. % or greater more preferably 2.1 wt % or greater asthis produces even higher creep resistance. Ruthenium is preferablylimited to 3.0 wt. % as this gives the preferred balance between costand creep resistance. A more preferred maximum level of ruthenium is 2.9wt % yet further to reduce cost, whilst still benefiting from a highcreep merit index. If the tungsten content is limited to 11.5 wt. % orless, the alloy density can be decreased to 9.0 g/cm³ or less.Preferably the tungsten content is limited to 9.5 wt. % as this producesan alloy with an even lower density (FIGS. 6 and 10). Lower levels oftungsten also ensure microstructural stability (FIGS. 11-14).

From FIGS. 7-10 a minimum content of rhenium of 3.5 wt. % or more isshown to produce a high creep merit index. Preferably the rheniumcontent is greater than 4.5 wt. % as this produces an alloy with abetter balance between density (FIG. 6) and creep resistance (FIG. 10).Even more preferable is an alloy containing at least 5.3 wt. % ofrhenium as this composition produces an alloy with an even betterbalance of creep resistance, density. In such an alloy cost can also bereduced as lower levels of ruthenium may be required (FIG. 9).

Molybdenum behaves in a similar way to tungsten i.e. this slow diffusingelement can improve creep resistance. Therefore, it is preferred thatmolybdenum is present in an amount of at least 0.1 wt %. However,additions of molybdenum must be controlled as it strongly increases thealloys propensity to form deleterious TCP phases. Therefore, molybdenumis limited to 0.5 wt. % or less.

From FIGS. 7-10 and a knowledge that molybdenum can substitute tungsten,it can be determined that a good level of creep resistance is achievedwhen additions of tungsten, rhenium, molybdenum adhere to the followingEquation,

ƒ(Creep)=2.92W _(Re)+(W _(W) +W _(Mo))

where, f(Creep) is a numerical value which is greater than or equal to21.9 and W_(Mo) is the weight percent of molybdenum in the alloy. Thisproduces an alloy with a creep merit index as calculated of 12×10⁻¹⁵ m⁻²s or more. Preferably the numerical value f(Creep) is greater than 24.6as this produces an alloy with increased creep resistance. Additionallythis allows lower levels of ruthenium to attain equivalent creepresistance thus reducing cost.

Additions of cobalt are optional. However, modelling calculations showthat cobalt increases the creep merit index. Additions of cobalt arealso know to lower the stacking fault energy in the gamma matrix whichalso improves creep resistance. Furthermore, additions of cobalt canimprove the ease of processing as it can lower the γ′ solvustemperature, helping to increase the solutioning window. Thus, at least0.1 wt % cobalt is desirably present. FIG. 15 shows that cobaltadditions lower the γ′ solvus temperature when the aluminium to tantalumratio is in the most preferred range 0.75-1.36. A preferred lower limitof cobalt is 7.0 wt. % as this produces an alloy with improved creepresistance and a lower γ′ solvus temperature which is beneficial forheat treatment processes. However, cobalt additions must be limited ashigh cobalt levels will increase the alloy's creep anisotropy,particularly in primary creep. This makes the creep rate stronglydependent upon orientation of the single crystal. An upper limit of 12.0wt. % cobalt controls the amount of creep anisotropy to an acceptablelevel. A preferred upper limit is 11.0 wt. % as creep anisotropy is evenless prevalent. A more desirable upper limit is 10.9 wt % even furtherto decrease the chance of creep anisotropy.

In order to remain resistant to creep over a significant time period theaddition of slow diffusing elements rhenium, tungsten and ruthenium isrequired. Additions of chromium are also required to promote resistanceto oxidation/corrosion damage. However, the addition of high levels oftungsten, rhenium and chromium were found to increase the propensity toform unwanted TCP phases, primarily σ, P and μ phases. FIGS. 11-14 showsthe effect of chromium, tungsten and rhenium additions on the overallfraction of TCP phases (σ+μ+P). Preferably the additions of theseelements were controlled to ensure that the levels of the TCP phaseswere equivalent to or lower than current fourth generation superalloys(Table 3).

The minimum chromium content for the present invention is greater thanor equal to 3.5 wt. % and preferably greater than or equal to 4.0 wt. %in order to attain oxidation resistance which is improved in comparisonto current fourth generation single crystal alloys which have Crcontents ranging between 2.0-3.2 wt. %. That is, a higher weight percentof chromium is provided than in the current fourth generation alloys onthe basis that this will improve oxidation resistance compared to thosealloys. The chromium content is limited to 6.5 wt. % to reduce thepropensity for the alloy to form the deleterious TCP phases (FIGS.11-14). Preferably the chromium content in the alloy is limited to 5.0wt. % as this produces an alloy with the best balance between oxidationresistance and microstructural stability. In the present invention therhenium content in the alloy is limited to 7.0 wt. % or less (to ensureacceptable microstructural stability, FIG. 14) and more preferably 6.0wt. % or less as rhenium at a level of between 4.5 wt. % and 6.0 wt. %provides a good balance between density, creep resistance andmicrostructural stability. Based upon the rhenium levels allowable for abalance between microstructural stability, density and creep resistancethe minimum tungsten level required for the present invention is 4.5 wt.% or more, as this provides a balance between creep resistance (FIGS.7-10), cost and microstructural stability (FIGS. 11-14). In order toachieve high creep resistance (FIGS. 7-10) a preferred minimum level oftungsten is 7.0 wt. %, desirably at least 7.1 wt %.

It is beneficial that when the alloy is produced, it is substantiallyfree from incidental impurities. These impurities may include theelements carbon (C), boron (B), sulphur (S), zirconium (Zr) andmanganese (Mn). If concentrations of carbon remain at 100 PPM or below(in terms of mass) the formation of unwanted carbide phases will notoccur. Boron content is desirably limited to 50 PPM or less (in terms ofmass) so that formation of unwanted boride phases will not occur.Carbide and boride phases tie up elements such as tungsten or tantalumwhich are added to provide strength to the γ and γ′ phases. Hence,mechanical properties including creep resistance are reduced if carbonand boron are present in greater amounts. The elements Sulphur (S) andZirconium (Zr) preferably remain below 30 and 500 PPM (in terms ofmass), respectively. Manganese (Mn) is an incidental impurity which ispreferably limited to 0.05 wt % (500 PPM in terms of mass). The presenceof Sulphur above 0.003 wt. % can lead to embrittlement of the alloy andsulphur also segregates to alloy/oxide interfaces formed duringoxidation. This segregation may lead to increased spallation ofprotective oxide scales. The levels of zirconium and manganese must becontrolled as these may create casting defects during the castingprocess, for example freckling. If the concentrations of theseincidental impurities exceed the specified levels, issues surroundproduct yield and deterioration of the material properties of the alloyis expected.

Additions of hafnium (Hf) of up to 0.5 wt. %, or more preferably up to0.2 wt. % are beneficial for tying up incidental impurities in thealloy, in particular carbon. Hafnium is a strong carbide former, soaddition of this element is beneficial as it will tie up any residualcarbon impurities which may be in the alloy. It can also provideadditional grain boundary strengthening, which is beneficial when lowangle boundaries are introduced in the alloy.

Additions of the so called ‘reactive-elements’, Silicon (Si),Yttrium(Y), Lanthanum (La) and Cerium (Ce) may be beneficial up tolevels of 0.1 wt. % to improve the adhesion of protective oxide layers,such as Al₂O₃. These reactive elements can ‘mop-up’ tramp elements, forexample sulphur, which segregates to the alloy oxide interface weakeningthe bond between oxide and substrate leading to oxide spallation. Inparticular, it has been shown that additions of silicon to nickel basedsuperalloys at levels up to 0.1 wt. % are beneficial for oxidationproperties. In particular silicon segregates to the alloy/oxideinterface and improves cohesion of the oxide to the substrate. Thisreduces spallation of the oxide, hence, improving oxidation resistance.

Based upon the description of the invention presented in this section,broad and preferred ranges for each elemental addition were defined,these ranges are listed in Table 4. An example composition—alloyABD-2—was selected from the preferred compositional range, thecomposition of this alloy is defined in Table 4. Alloy ABD-2 was foundto be amenable to standard methods used for the production of singlecrystal turbine blade components. This production method involves:preparation of an alloy with the composition of ABD-2, preparation of amould for casting the alloy using investment casting methods, castingthe alloy using directional solidification techniques where a ‘grainselector’ is used to produce a single crystal alloy, subsequentmulti-step heat treatment of the single crystal casting.

TABLE 4 Compositional range in wt. % for the newly design alloy. BroadPreferred Nominal Min Max Min Max ABD-2 Cr 3.5 6.5 4.0 5.0 4.0 Co 0.012.0 7.0 11.0 9.0 W 4.5 11.5 7.0 9.5 7.4 Al 3.7 6.8 5.1 6.6 6.4 Ta 5.09.0 5.0 7.3 5.6 Mo 0.0 0.5 0.1 0.5 0.0 Re 3.5 7.0 4.5 6.0 5.6 Ru 1.0 3.72.0 3.0 2.6 Hf 0.0 0.5 0.0 0.2 0.0 Nb 0.0 0.5 0.0 0.5 0.0 Ti 0.0 0.5 0.00.5 0.0 V 0.0 0.5 0.0 0.5 0.0 Si 0.0 0.1 0.0 0.1 0.0 Ce 0.0 0.1 0.0 0.10.0 Y 0.0 0.1 0.0 0.1 0.0 La 0.0 0.1 0.0 0.1 0.0

Experiment testing of alloy ABD-2 was used to validate the key materialproperties aimed at with the alloy of the invention, mainly sufficientcreep resistance and improved oxidation behaviour in comparison to thatof a current single crystal alloys used for IGT applications. Thebehaviour of alloy ABD-2 was compared with alloy TMS-138-A, which wastested under the same experimental conditions.

Single crystal castings of alloy ABD-2 of nominal composition accordingto Table 4 were manufactured using conventional methods for producingsingle crystal components. The castings were in the form of cylindricalbars of 10 mm diameter and 160 mm in length. The cast bars wereconfirmed to be single crystals with an orientation within 10° from the<001> direction.

The as cast material was given a series of subsequent heat treatments inorder to produce the required γ/γ′ microstructure. A solution heattreatment was conducted at 1325° C. for 6 hours, this was found toremove residual microsegregation and eutectic mixtures. The heattreatment window for the alloy was found to be sufficient to avoidincipient melting during the solution heat treatment. Following thesolution heat treatment the alloy was given a two stage ageing heattreatment, the first stage conducted at 1120° C. for 2 hours and thesecond stage conducted at 870° C. for 16 hours.

Creep specimens of 20 mm gauge length and 4 mm diameter were machinedfrom fully heat-treated single crystal bars. The orientation of the testspecimens were within 10° from the <001> direction. Test temperaturesranging from 800 to 1100° C. were used to evaluate the creep performanceof the ABD-2 alloy. Cyclic oxidation tests were performed on the fullyheat treated material. Cyclic oxidation tests were carried out at 1000°C. using 2 hours cycles over a time period of 50 hours.

A Larson-Miller diagram was used to compare the creep resistance ofalloy ABD-2 with alloy TMS-138A. In FIG. 16 a comparison of time to 1%creep strain is presented for both alloys. The time to 1% strain iscritical as most gas turbine components are manufactured to tighttolerances to achieve maximum engine performance. After low levels ofstrain—in the order of a few percent—components will often be replaced.It is seen that alloy ABD-2 is comparable to TMS-138A in time to 1%creep strain. FIG. 17 shows a comparison of time to creep rupture forboth alloys, it is seen that alloy ABD-2 has a rupture life comparableto that of TMS-138A.

The oxidation behaviour of alloys ABD-2 and TMS-138A was also compared.As turbine temperatures continue to rise—improving thermal efficiency ofthe engine—component failure due to corrosion damage such as oxidationis becoming more prevalent. Hence, significant gains in component lifemay be attained by improving oxidation/corrosion resistance. The alloyABD-2 was designed such that it would have improved oxidation behaviourrelative to current second generation alloys. Cyclic oxidation resultsfor ABD-2 and TMS-138A are presented in FIG. 18. A reduction in massgain with respect to time is evidence of improved oxidation behaviour asthe formation of a protective oxide scale has occurred limiting theingress of oxygen into the substrate material. The ABD-2 alloy showssignificantly reduced weight gain with respect to time when compared toTMS-138A, indicative of improved oxidation performance.

Overall the alloy ABD-2 shows equivalent creep behaviour in comparisonto TMS-138A. This has been achieved using an alloy with a significantlyimproved oxidation behaviour. Thus, design goals have been met whilststill achieving a low cost and density alloy which is amenable toconventional manufacturing techniques.

1. A nickel-based alloy composition consisting, in weight percent, of:between 3.5 and 6.5% chromium, between 0.0 and 12.0% cobalt, between 4.5and 11.5% tungsten, between 0.0 and 0.5% molybdenum, between 3.5 and7.0% rhenium, between 1.0 and 3.7% ruthenium, between 3.7 and 6.8%aluminium, between 5.0 and 9.0% tantalum, between 0.0 and 0.5% hafnium,between 0.0 and 0.5% niobium, between 0.0 and 0.5% titanium, between 0.0and 0.5% vanadium, between 0.0 and 0.1% silicon, between 0.0 and 0.1%yttrium, between 0.0 and 0.1% lanthanum, between 0.0 and 0.1% cerium,between 0.0 and 0.003% sulphur, between 0.0 and 0.05% manganese, between0.0 and 0.05% zirconium, between 0.0 and 0.005% boron, between 0.0 and0.01% carbon, the balance being nickel and incidental impurities.
 2. Thenickel-based alloy composition according to claim 1, consisting, inweight percent, of between 4.0 and 5.0% chromium.
 3. The nickel-basedalloy composition according to claim 1, consisting, in weight percent,of between 0.1 to 12.0 wt % cobalt, more preferably 0.1 to 11.0% cobalt,or 0.1 to 10.9% cobalt, most preferably 7.0 and 11.0% cobalt or evenbetween 7.0 and 10.9% cobalt.
 4. The nickel-based alloy compositionaccording to claim 1, consisting, in weight percent, of between 7.0 and9.5% tungsten, or of between 7.1 and 11.5 wt % tungsten, preferablybetween 7.1 and 9.5% tungsten.
 5. The nickel-based alloy compositionaccording to claim 1, consisting, in weight percent of, between 3.7 and6.6%, preferably between 5.1 and 6.6% aluminium, more preferably between5.5 and 6.6% aluminium.
 6. The nickel-based alloy composition accordingto claim 1, consisting, in weight percent, of between 5.0 and 7.3%tantalum.
 7. The nickel-based alloy according to claim 1, consisting, inweight percent, of at least 0.1% molybdenum.
 8. The nickel-based alloyaccording to claim 1, consisting, in weight percent, of between 4.5 and6.0% rhenium, more preferably between 5.3 and 6.0% rhenium.
 9. Thenickel-based alloy according to claim 1, consisting, in weight percent,of between 2.0 and 3.0% ruthenium, preferably between 2.1 and 2.9%ruthenium.
 10. The nickel-based alloy composition according to claim 1,wherein the following equation is satisfied in which W_(Ta) and W_(Al)are the weight percent of tantalum and aluminium in the alloyrespectively33≤W _(Ta)+5.1W _(Al)≤39.
 11. The nickel-based alloy compositionaccording to claim 1, wherein the following equation is satisfied inwhich W_(Ta) and W_(Al) are the weight percent of tantalum and aluminiumin the alloy respectively2.2≤5.15W _(Al)−0.5W _(Al) ² −W _(Ta),preferably 2.9≤5.15W _(Al)−0.5W _(Al) ² −W _(Ta).
 12. The nickel-basedalloy composition according to claim 1, wherein the following equationis satisfied in which W_(Ru) and W_(Re) are the weight percent ofruthenium and rhenium in the alloy respectively4.5≥W _(Ru)+0.225W _(Re),preferably 3.9≥W _(Ru)+0.225W _(Re).
 13. The nickel-based alloycomposition according to claim 1, wherein the following equation issatisfied in which W_(Re) and W_(W) are the weight percent of rheniumand tungsten in the alloy respectively15.8≥1.13W _(Re) +W _(W),preferably 14.4≥1.13W _(Re) +W _(W).
 14. The nickel-based alloycomposition according to claim 1, wherein the following equation issatisfied in which W_(Re), W_(Mo) and W_(W) are the weight percent ofrhenium, molybdenum and tungsten in the alloy respectively21.9≤2.92W _(Re)+(W _(W) +W _(Mo)),preferably 24.6≤2.92W _(Re)+(W _(W) +W _(Mo)).
 15. The nickel-basedalloy composition according to claim 1, wherein the sum of the elementsniobium, titanium and vanadium, in weight percent, is less than 1%,preferably less than 0.5 wt %.
 16. The nickel-based alloy compositionaccording to claim 1, consisting, in weight percent, of between 0.0 and0.2% hafnium.
 17. The nickel-based alloy composition according to claim1, having between 60% and 70% volume fraction γ′.
 18. The nickel-basedalloy according to claim 1, wherein the sum of the elements niobium,titanium, vanadium and tantalum, in weight percent, is between 5.0 and9.0%, preferably between 5.0-7.3%.
 19. A single crystal article formedof the nickel-based alloy composition of claim
 1. 20. A turbine bladefor a gas turbine engine formed of an alloy according to claim
 1. 21-25.(canceled)